Processes for producing thicker gage products of niobium microalloyed steel

ABSTRACT

A process for controlling austenite grain size in austenite processing through nano-scale precipitate engineering of TiN—NbC composites to produce thicker gage product of niobium microalloyed line pipe steel that includes controlling the base chemistry of the steel to include 0.003-0.004 wt. % nitrogen, 0.012-0.015 wt. % titanium, 0.03-0.07 wt. % carbon, and 0.07-0.15 wt. % niobium; conducting a first stage of roughing above the temperature of dissolution of NbC to refine austenite grain size by static recrystallization below 30 microns, cooling so that the center of the thick rolled slab is about 1040° C., conducting a second stage of roughing to promote strain induced growth of NbC on pre-existing TiN to form TiN—NbC composites to increase Zener pinning pressure on austenite grain boundaries to about 0.1 MPa in order to prevent recrystallization and grain coarsening of austenite grains beyond 30 microns at the end of roughing, and conducting finish rolling below 920° C.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a continuation-in-part of U.S. patent application Ser. No. 14/325,940 filed on Jul. 8, 2014, and published as US 2016/0010190 A1 on Jan. 1, 2016, the entire contents of which are incorporated herein by reference.

BACKGROUND OF THE INVENTION 1. Field of the Invention

This invention relates to austenite grain size control by preventing grain coarsening of austenite in upstream processing of niobium microalloyed steel in order to produce thicker gage products with excellent drop weight tear test (DWTT) toughness (as measured in accordance with API RP 5L3 (Apr. 1, 1996)). In current technological practice, there is no intentional control measure to prevent grain coarsening of austenite before entry to finish rolling. As a consequence, heavy rolling reductions are applied to coarse grained austenite during finish rolling in order to increase the surface to volume ratio by geometric means through heavily pancaked austenite. Heavy rolling reductions applied in the finishing mill, often approaching limits of mill loading, inevitably limit thickness of the final product. In order to improve the safety and efficiency of the transport of natural gas and oil through the pipe lines, there is a growing demand for thicker gage pipes, particularly for deep offshore projects.

This invention targets austenite grain size control upstream i.e., at high temperatures in order to produce thicker gage product. This invention utilizes the formation of nano-scale TiN—NbC composite precipitates to pin austenite grain boundaries, and prevent them from coarsening at high temperatures (>980° C.) so that less rolling reduction and pancaking of austenite is required to obtain target properties of high strength and excellent toughness at low temperature as measured by ductile to brittle transition temperature and percentage shear in drop weight tear tests. There are additional benefits related to reduced texture related anisotropy of properties and improved ductile fracture arrestability of pipes in service resulting from engineering TiN—NbC composites by controlling the spacing of the precipitates by TiN independently from the size of precipitates by strain induced growth of NbC on pre-existing TiN to control the Zener pinning pressure of precipitates on austenite grain boundaries.

2. Description of Related Art

Although austenite grain size at the end of roughing is fine (<30 microns), significant grain coarsening of austenite occurs subsequent to the end of roughing, which must be prevented from occurring in order to produce a thicker gage product. If the austenite grain size is coarse, it is feasible to apply heavy rolling reductions to “pancake” the austenite and increase the surface to volume ratio of austenite, which increases the nucleation sites for ferrite upon transformation. To the extent that grain coarsening is controlled by a diffusion mechanism which is dependent on time and temperature, rapid thermal cooling should decrease the kinetics of diffusion for grain coarsening. However, it has not been found feasible to adequately cool the center of thick transfer bars to prevent grain coarsening. Nevertheless, accelerated cooling upstream is still beneficial to avoid rolling in the partial recrystallization regime and avoid depletion of solute niobium by excessive growth of niobium carbide precipitates upstream. Therefore, there is a need to develop alternative strategies to prevent grain coarsening. Although solute niobium retards grain coarsening by retarding boundary mobility, the magnitude of solute drag on boundary mobility is weak at high temperatures. As a consequence, even though most of the niobium is available as solute in conventional processing, it is not found to be effective in preventing grain coarsening in the described conditions, e.g. where the temperature is greater than 1100° C. during roughing. As shown in FIGS. 15A-C, traditional static recrystallization results in grain coarsening of fine austenite grains. In the prior art described by U.S. Patent Publication No. 2003/0190251 to Bai et al., austenite grains of less than 30 microns are obtained at the end of roughing in a Steckel mill for a transfer bar of 31 mm thickness. After a short holding time of 80 seconds, the transfer bar is transported to finish mill for finish rolling to produce a coil of 9.1 mm thickness. However, even during this short holding period, austenite grains will coarsen to 60 microns, which is not recognized in the prior art. The smaller the radius of the austenite grains, the larger the driving force is for grain coarsening by capillarity.

In a modern Steckel mill engaged in producing 19 mm gage thickness, the cooling time between the end of roughing and the start of finish rolling is typically about 300 seconds. Thus, grain coarsening of austenite during the holding period is a significant problem. FIG. 15A shows morphology of prior-austenite grains after 25% deforming and holding for 3 seconds at 1000° C. FIG. 15B shows morphology of prior-austenite grains after 25% deforming and holding for 60 seconds at 1000° C. FIG. 15C shows morphology of the prior-austenite grains after 25% deforming and holding for 240 seconds at 1000° C. Additional information regarding the grain coarsening found in traditional austenite grains can be found in C. L. Miao et al., “Recrystallization and strain accumulation behaviours of high Nb-bearing line pipe steel in plate and strip rolling” Mat. Sci. and Engineering A, 527 (2010), 4985-4992. A quantitative analysis of this problem shows that austenite grains of less than 30 microns obtained at the end of roughing will coarsen to more than 80 microns, see FIG. 16. As shown in FIG. 16, significant grain coarsening of refined austenite grains occurs (for traditional processes) from the end of rough rolling to the start of finish rolling. In FIG. 16, the linear segment indicates temperature and the curved segment indicates diameter. In 300 seconds, which is what the bar takes for transfer, the grains coarsen to 85 microns. In this case, solute Nb is assumed to be 0.04 wt %.

One of the aims of embodiments of the present invention is to prevent grain coarsening between the end of roughing and the start of finish rolling. Thus, an object of embodiments of the present invention is to prevent grain coarsening of austenite grains by metallurgical means, through pinning austenite grain boundaries by second phase particles using the Zener pinning mechanism. The use of TiN particles to pin austenite grain boundaries by the Zener pinning mechanism is well established and has been disclosed in prior patents (see, for example, U.S. Pat. No. 6,899,773; U.S. Pat. No. 6,183,573; and U.S. Pat. No. 5,900,075). While these patents identify conditions under which high number density of TiN precipitates can be promoted, the limiting austenite size achievable by TiN alone is typically 60 to 80 microns in the high temperature window of processing. NbC is sluggish to nucleate by itself and is aided by dislocations generated by deformation to promote strain induced nucleation of NbC, which is associated with large undercooling. Strain induced precipitation of NbC is used in controlled rolling of microalloying technology, where strain induced precipitation of NbC is used to pin austenite grain boundaries during thermo-mechanical controlled rolling during the low temperature window of processing. However, by promoting epitaxial growth of NbC on pre-existing TiN in accordance with embodiments of the present invention, TiN—NbC composite precipitates are obtained with negligible undercooling in the high temperature window towards the end of roughing. These nano-scale TiN—NbC composite precipitates are available to pin austenite grains towards the end of roughing and limit austenite grain size to under about 30 microns on entry to finish rolling, which is essential to produce thicker gage in line pipe grades.

Accordingly, it is an object of the invention to reduce the need for large rolling reductions and heavy pancaking during finish rolling in order to obtain increased gages of finished product.

It is another object of the invention to produce uniform fine austenite grain size before pancaking and apply less pancaking to produce thicker gage product, which exhibits homogeneous properties without anisotropy due to unfavorable crystallographic texture development.

It is yet another object of the invention to obtain consistently low ductile to brittle transition temperature (DBTT) and good drop weight tear test (DWTT) performance. DWTT properties are empirically correlated with thickness of pancaked austenite grain. By refining the austenite grain size, the thickness of pancaked austenite grains can be decreased to meet target DWTT properties.

SUMMARY OF THE INVENTION

It has now been discovered that the addition of niobium to titanium-bearing super-martensitic stainless steel refines the austenite grain size due to the formation of titanium-niobium bearing composite precipitates. This led to the present invention's development of nano-scale precipitation engineering of TiN—NbC composite precipitates to prevent austenite grains from coarsening. According to embodiments of the present invention, TiN precipitates, which are formed just after solidification in the continuous cast slab, are used to control the interparticle spacing, while NbC precipitates growing on pre-existing TiN particles are used to control the size of the precipitates, both size and spacing of TiN—NbC composite precipitates are the key to pinning austenite grains of the required size to prevent them from coarsening. The driving force for grain coarsening is capillary force, which can be determined from the equation: capillary force=2 γ/R, where R is the radius of curvature of the grain boundary and γ is the surface energy of the boundary.

The driving force for grain coarsening tends to decrease as the grain size increases. In accordance with embodiments of the present invention, a 30 micron grain size is targeted instead of the conventional 60 microns. This driving force for boundary movement is counteracted if particles pin the boundary. The pinning force increases with the number density and size of the particles. Thus, the driving force for grain coarsening when the target grain size is 30 microns can be determined from the number density of particles [TiN], which sets up the interparticle spacing that can be measured, e.g., 200 nm. But the particle size of TiN is too small, about 15 nm. The limiting austenite grain size is 90 microns, which is rather coarse. By growing NbC, TiN—NbC composites can be formed, which are now large, about 25 nm. The limiting austenite grain size is 32 microns, which is close to target. By growing to 30 nm, as can be seen from Table 1 below, the limiting austenite grain size is decreased to 22 microns. It should be noted that the number density and the volume fraction of precipitates are controlled by the thermodynamics and kinetics of precipitation which, in turn, depend upon the chemical composition and processing parameters of the steel.

TABLE 1 Increasing N Increasing base Nb concentration concentration Particle Diameter, nm Distance, nm 10 15 20 25 30 35 40 50 60 70 80 150 85 38 21 14 10 7 5 3 2 1.8 1.3 200 203 90 51 32 22 17 13 8 6 4 3 250 397 176 99 63 44 32 25 16 11 8 6 300 687 305 171 110 76 56 43 28 19 14 11 350 1091 485 273 174 121 89 68 44 30 22 17 400 1629 724 407 260 181 133 102 65 45 33 25 450 2320 1030 580 371 257 190 145 93 65 47 36 500 3183 1414 795 509 353 259 199 127 88 65 50 550 4236 1882 1059 678 471 346 265 170 118 86 66 Zener limiting Austenite grain size in micrometers

The validity of the mechanism underpinning the technology of nano-scale precipitation engineering for austenite grain size control in upstream processing is demonstrated by experimental results on line pipe grades processed under plate rolling and hot strip rolling conditions.

In conventional processes, roughing is carried out in a high temperature window where there is no significant thermodynamic potential for precipitation of NbC and Zener drag from TiN precipitates is inadequate to stop recrystallization. Austenite grains are refined in roughing by static recrystallization that follows each pass, and austenite grains are refined below 30 microns at the end of roughing. However, in conventional processes, austenite grains tend to coarsen during the cooling time before entry to finish rolling temperature.

In traditional processes, there is no deliberate step to prevent grain coarsening of austenite to 60 microns. According to the embodiments of present invention, there is a second stage of roughing, comprising one or two heavy passes towards the end of roughing after cooling the center of thick sections to a low temperature, where there is adequate thermodynamic potential for strain induced growth of NbC on pre-existing TiN to form TiN—NbC. This second stage of roughing is conducted when the center of the steel product is at about 1040° C., well below the dissolution temperature of NbC, so that NbC can grow epitaxially on pre-existing TiN upon deformation to promote strain induced growth of NbC on pre-existing TiN. This increases the Zener pinning pressure, e.g. drag, exerted by the TiN—NbC precipitates on austenite grain boundaries to counteract the driving force for recrystallization, and arrest recrystallization. Since the driving force for grain coarsening of 30 microns is less than that for recrystallization, austenite grain coarsening is limited to a maximum of 30 microns. Thus, temperature of no recrystallization (Tnr) is reached in the center of the transfer bars at the end of the second stage of roughing, which, as described below, acts to suppress rotated cube texture and prevents grain coarsening in the long cooling time between the end of roughing and the start of finish rolling.

The cubic texture of recrystallized austenite grains upon transformation promotes undesirable rotated cube texture, which promotes brittle cleavage fracture. In embodiments of the present invention, by work-hardening the recrystallized austenite grains, the texture is changed to brass or copper texture, which upon transformation gives desirable texture, shown in the bottom half of FIG. 17, thereby suppressing the undesirable rotated cube texture. The top half of FIG. 17 illustrates that undesirable rotated cube texture will occur in the center of thicker gage product (16-32 mm gage thickness), which is at a temperature higher than the temperature of no recrystallization by strain induced precipitation of NbC as in conventional processing. Since the temperature in the center of thick transfer bars (60 mm) is typically about 100° C. higher than the surface, it is essential to raise the Tnr correspondingly by about 100° C. in order to pancake the center of thick sections. Since nano-scale precipitate engineering of TiN—NbC composite raises Tnr by about 100° C. compared with conventional processing by strain induced precipitation of NbC, it is another object of the present invention to suppress undesirable rotated cube texture in thicker gage that promotes cleavage fracture.

By controlling the spacing and size of TiN—NbC, delayed strain induced precipitation of NbC is suppressed, thereby retaining adequate solute niobium in the matrix to suppress polygonal ferrite nucleation. In the prior art by Bai et al., the final microstructure is described as 30% polygonal ferrite and 70% acicular ferrite. It is the objective of the current invention to suppress undesirable delayed strain induced precipitation of NbC in austenite that promotes polygonal ferrite in the final microstructure. It is the objective of the embodiments of the present invention to promote acicular ferrite transformed at low temperature in order to obtain high number density of high angle boundaries that suppress brittle fracture.

This also acts to suppress rotated cube texture and prevents grain coarsening in the long cooling time between the end of roughing and the start of finish rolling. In the case of processing heavier gage in a conventional hot strip mills to produce coil, one heavy pass after cooling the center to Tnr by nano-scale precipitate engineering is found to be adequate for a gage thickness of 20 mm, see FIG. 18, which shows a process in accordance with embodiments of the present invention.

Embodiments of the present invention provide a platform for austenite grain size control in upstream processing of austenite of niobium microalloyed steels to which downstream processing and final properties of the product are related.

A process for controlling austenite grain size in austenite processing of niobium microalloyed steel through nano-scale precipitation engineering of TiN—NbC composites to produce thicker gage product comprises controlling the base chemical composition of a steel product to include about 0.003-0.004 wt % nitrogen, 0.012-0.015 wt % titanium, 0.03-0.07 wt % carbon, and 0.07-0.15 wt % niobium; lowering the temperature of roughening to end the roughening operation in the temperature range of from about 940° C. to 1040° C., e.g. 960° C. at the surface; retaining greater than about 0.03 wt % niobium in solution in the matrix by rapid cooling of the product to enter the finish rolling operation below the temperature of no recrystallization, with an austenite grain size of about 30 microns; and applying reduced rolling reduction in the finish rolling operation. The temperature for the last roughening/roughing pass is usually based on surface temperature, but can also be based on the temperature at the center of the thickest section of the steel product, which is about 100° C. higher than the surface temperature. For example, the temperature for the last roughing deformation pass can be defined as 1040° C. in the center of the rolled section, but in the mill, it has to be inter-related to surface temperature. Lowering the temperature of roughening prevents grain refined austenite from coarsening above about 30 microns by formation of TiN—NbC composite precipitates. Applying reduced rolling reduction in the finish rolling operation acts to pancake the fine austenite grain size of about 30 microns to obtain a sufficient surface to volume ratio to produce thicker gage resulting steel product.

In accordance with some embodiments, a process for controlling austenite grain size in austenite processing through nano-scale precipitation engineering of TiN—NbC composites to produce thicker gage product of niobium microalloyed steel for line pipe application with good crack arrestability to ductile fracture propagation includes controlling the base chemical composition of a steel product to include about 0.003-0.004 wt % nitrogen, 0.012-0.015 wt % titanium (to control the interparticle spacing of TiN to be about 200 nm), 0.03-0.07 wt % carbon, and 0.07-0.15 wt % niobium (to control thermodynamic potential for precipitation of NbC). The process includes carrying out the following process steps in roughing of the steel product (e.g. a reheated slab) with austenite grain size of 200-300 microns and the chemistry designed to give TiN precipitates of 10-15 nm size with an interparticle spacing of 200-300 nm.

Process Step-1: Conducting the first stage of roughing passes in high temperature window (about 1080° C. to 1150° C.), in which there is no significant thermodynamic potential for precipitation of NbC, as found in traditional processes, in order to refine the austenite grains from 300 microns to less than 30 microns by repeated static recrystallization.

Process Step-2: Holding for 100-120 seconds for temperature homogenization and the center of the section to cool to a temperature typically about 1040° C. at which there is significant thermodynamic potential for precipitation of NbC to occur. The surface temperature of the thick section of the steel product (slab) will be significantly less than the center, by up to 100° C. less. This will increase thermodynamic potential for precipitation of NbC available in the surface than in the center.

Process Step-3: Conducting a second stage of roughing, e.g. rough rolling, that includes applying one or two heavy pass reductions (>25% reduction relative to steel slab after the first stage of roughing) when the center of the section is at about 1040° C., in order to promote strain induced growth of NbC on pre-existing TiN. This will be aided by accelerated diffusion by dislocations generated by deformation to form TiN—NbC composites of 25-30 nm with an interparticle spacing of 200-300 nm. These precipitates will exert Zener pinning pressure of 0.1 MPa to stop recrystallization in the center of thick section, thereby promoting Tnr at about 1040° C. This enables the center of thick transfer bars (e.g. the steel product after the second stage of rough rolling) at high temperature to be fully pancaked to suppress the occurrence of undesirable rotated cube texture from occurring in the center of thicker gage products. The Zener pinning pressure of 0.1 MPa is more than adequate to counteract the driving force for grain coarsening by capillarity of austenite grains of about 30 microns size.

Process step-4: Conducting rapid cooling to enter finish rolling so that the transfer bar has a center temperature below 920° C. so that adequate solute niobium (Nb>0.03 wt %) is retained in the matrix for transformation hardening. The typical cooling rate ranges from 1 to 0.5° C./second.

Process Step-5: Applying an adequate rolling reduction of about 70% to pancake the austenite grains of about 30 microns in finish rolling to obtain large surface to volume ratio (Sv factor) of the austenite grains with large strain accumulation and so that there is increased surface area of austenite that acts as nucleation sites for ferrite as well as high dislocation density to increase the number density of nucleation sites per unit area of austenite grain surface. The aim is to control thickness of pancaked austenite grains to be less than 10 microns. In the absence of Zener pinning pressure from TiN—NbC, austenite will coarsen to 60 microns, as exemplified by the conventional practices used in the prior art.

Process Step-6: Conducting accelerated cooling of austenite grains with large Sv factor, large strain accumulation and adequate solute niobium (0.03 wt % Nb) in the matrix to promote transformation of acicular ferrite at low temperature, e.g. laminar water cooling at a rate ranging from 8-15° C./s, or higher in modern mills (10-30° C./s), and to promote acicular ferrite grains of 1-3 microns with high density and dispersion of high angle boundaries to suppress brittle fracture to produce thicker gage with enhanced crack arrestability to ductile fracture propagation and/or in order to promote crystallographic variants that give high density and dispersion of high angle boundaries to suppress brittle fracture. Those skilled in the art will readily appreciate that ferrite grains nucleate on austenite grain surfaces and therefore, by pancaking the austenite, we increase the surface area per unit volume of austenite. This is a geometric parameter that increases the nucleation sites for ferrite. The kinetics of nucleation is aided by dislocation density, which increases the number density of acicular ferrite nucleating per unit area of austenite boundary (Rate of nucleation per unit area of austenite surface—Kinetics). I. Kozazu, G. Ouchi, T. Sampei and T. Okita, Proc. of International conf. on Microalloying 75, Union Carbide, New York, (1976), 100.

The increased Tnr by nano-scale precipitate engineering of TiN—NbC composite to 1040° C. instead of 940° C. for Tnr by NbC is adequate to pancake austenite grains occurring in the center of thicker sections before the transformation, thereby suppressing undesirable rotated cube cleavage texture from occurring in the center of thicker gage steel product. The control of interparticle spacing of TiN—NbC composite is effective in suppressing delayed strain induced precipitation of NbC. In conventional practice, delayed strain induced precipitation of NbC occurs, which renders it difficult to control residual solute niobium required to suppress polygonal ferrite formation in transformation product. Hardenability, in embodiments of the present disclosure, is controlled by adequate solute niobium to promote acicular ferrite transformation at low temperature by shear transformation to give high density and dispersion of high angle boundaries that suppress brittle fracture, thereby increasing percentage shear area in DWTT at low temperature. The control of interparticle spacing of TiN—NbC is effective in promoting austenite grain size to about 30 microns before pancaking thereby increasing the surface to volume ratio of austenite grains (Sv factor) that enables production of thicker gage.

The grain size can be controlled in the range of about 20-40 microns at entry to the finish rolling operation. TiN precipitates can be in the range of about 10-20 nm and the interparticle spacing can be about 200-300 nm. Thermodynamic potential for precipitation of NbC can occur towards the end of the roughing operation at temperatures ranging from about 940° C. to about 1040° C. TiN—NbC composites can be in the size range of about 20-50 nm. Accelerated cooling of the product can be applied to avoid rolling in the partial recrystallization regime. The process can include controlling nitrogen at or below about 40 ppm and making a titanium addition to meet the stoichiometric requirement to combine with all nitrogen to form high number density of TiN precipitate in about the 10-20 nm size range.

The process can include processing the steel product by conventional plate rolling, conventional hot strip rolling, steckel mill rolling, and/or near net shape processing. The steel product can be line pipe steel, infra-structure steel, and/or super-martensitic stainless steel. The crystallographic texture-related anisotropic properties of the resulting steel product can be minimized. The process can include substituting titanium partially or fully in the base chemistry with a member of the group consisting of Zr, Hf, Ta, W, V, Cr, Mo, Al and mixtures thereof, each with high affinity for nitrogen to form nano-scale precipitates on which NbC can grow epitaxially to give composite precipitates.

The process also can include partially substituting niobium in the base chemistry (with a minimum amount of 0.06 wt % niobium based on the available data) with other microalloying elements with high affinity for carbon selected from the group consisting of Zr, Hf, Ta, W, V, Cr, Mo and mixtures thereof, each to give composite precipitates. The process also can include substituting solute niobium on entry to finish rolling with other elements, which exhibit solute drag comparable to niobium. Still further, the process can include rapidly cooling the steel product to enter finish rolling with a center temperature at or below about 920° C. The steel product can exhibit a gage thickness of about 17-40 mm.

BRIEF DESCRIPTION OF THE DRAWINGS

FIGS. 1A and 1B show Electron Back-Scattered Diffraction (EBSD) Images, revealing austenite grain size of specimens of two 13% Cr-5% Ni-2% Mo supermartensitic stainless steels. FIG. 1A corresponds to the control steel without niobium additions, and FIG. 1B is the sample with 0.1 wt percent niobium addition, both steels contain titanium, as shown in Table 2. The addition of 0.1 wt percent niobium decreased the austenite grain size from 80 to 35 microns. i.e., titanium by itself could only produce an austenite grain size of 80 microns. But it is only with the addition of 0.1 wt. percent niobium, that austenite grain size can be decreased from 80 to 35 microns under identical processing conditions.

FIG. 2 is a TEM image of precipitates extracted on a carbon replica, showing nano-scale Ti—Nb bearing precipitates in the size range of 25-30 nm with a mean interparticle spacing of 230 nm. Energy dispersive analysis of the precipitates is shown alongside. The spectrum shows X-ray signals characteristic of titanium and niobium. The precipitates appear to be TiN—NbC composites similar to those found in line pipe steels, see FIGS. 4, 12 and 13.

FIG. 3 is a plot of driving force for grain coarsening of austenite as a function of austenite grain size. The coarsening is driven by reduction in surface energy of the grains. This is counteracted by particle pinning the boundary. The pinning pressure is governed by the particle size and number density. The number density determines interparticle spacing. Thus, small interparticle spacing and increased particle size are required to increase pinning pressure to counteract and prevent grain coarsening of fine grains. The particle limited grain size with and without niobium shows the effectiveness of TiN—NbC composites compared with TiN precipitates in pinning fine grains.

FIG. 4 is a high resolution TEM image of TiN—NbC composites obtained in low nitrogen line pipe steel microalloyed with titanium and niobium. Energy dispersive analysis shows X-ray signals characteristic of niobium and titanium in the composite precipitates. NbC precipitates appear to have grown on preexisting cuboidal TiN precipitates.

FIGS. 5A and 5B are salient results from prior work on the microstructural evolution of TiN—NbC composites in low interstitial titanium-niobium microalloyed steels investigated by hot torsion simulation of rolling. See, e.g., S. V. Subramanian, F. Boratto, J. J. Jonas and C. M. Sellars and published in the Proceedings of International Symposium on “Microalloyed Bar and Forging Steels” edited by Mike Finn, CIM, held at Hamilton, Ontario, Canada, Aug. 26-29, 1990, pp. 120-136. Based on quantitative analysis of thermodynamic potential for precipitation, mole fraction of TiN—NbC is plotted as a function of temperature. FIG. 5A shows the precipitate evolution curve for the high niobium low interstitial steel-G, containing carbon 0.03, nitrogen 0.003, titanium 0.014 and niobium 0.095 wt percent. Thermodynamic potential for precipitation of NbC starts at 1060° C. FIG. 5B shows the mean flow stress from hot torsion simulation (shown as open circles) as a function of the inverse of the absolute pass temperature for Steel-G. The bold line is the flow stress pertaining to a fully recrystallized steel. The onset of recrystallization retardation starts at 1060° C. corresponding to the onset of the thermodynamic potential for precipitation of NbC. Growth of NbC on preexisting precipitates of TiN is confirmed in this work, which obviates the need for independent nucleation of NbC. Thus, the resulting TiN—NbC composite retards recrystallization, causing the increase in flow stress detected by hot torsion rolling simulation results.

FIG. 6 is a schematic diagram that inter-relates the increase in size of TiN—NbC composite to volume fraction of NbC, which is determined by the thermodynamic potential for precipitation of NbC. The interparticle spacing is fixed by TiN on which NbC grows. This diagram illustrates that the rough rolling temperature window has to be lowered so that thermodynamic potential for growth of NbC is obtained on pre-existing TiN precipitates to form TiN—NbC composites at the end of rough rolling.

FIG. 7 is a process flow diagram of prior technology in which there is no intentional control of austenite grain size in upstream processing of rough rolling, and the austenite grain size on entry to finish rolling may be coarse, generally ranging in size from 60-80 microns. Therefore, heavy rolling reduction is applied in finish rolling stands downstream to reduce the thickness of pancaked austenite in order to obtain good toughness at low temperature in the final product. This limits the thickness of the final product generally well below 16 mm, processed by conventional plate rolling or conventional hot strip rolling of niobium microalloyed steel. This is illustrated with the specific example of Steel-A of 10 mm gage, with a high nitrogen content of 75 ppm. Rough rolling is carried out in the temperature window above the equilibrium temperature for precipitation of NbC. TEM characterization shows coarse precipitate of mean size 83 nm with a large interparticle spacing of 550 nm, which gives a Zener limiting austenite grain size of 62 microns. This requires heavy rolling reduction for pancaking austenite grains, resulting in thinner gage (<16 mm).

FIG. 8 is a process flow diagram based on the present invention wherein austenite grain size upstream is controlled by the size and spacing of TiN—NbC composite precipitates, which is referred to herein as “nano-scale precipitation engineering” or “nano-scale precipitate engineering.” The austenite grain size is intentionally controlled to be fine with a target grain size under 30 microns. This requires less rolling reduction to reduce the thickness of pancaked austenite grain size in order to obtain good toughness at low temperature as measured by percentage shear area in DWTT tests. As a result of applying less rolling reduction to the transfer bar, the thickness of the final product processed by conventional plate rolling or conventional hot strip rolling of niobium microalloyed steel can be increased well above 16 mm. This is demonstrated with the specific example of Steel-C. The steel contains a low nitrogen content of 0.004 wt percent and stoichiometric addition of Ti to combine with nitrogen. TEM characterization shows high number density of TiN precipitates with an interspacing of 220 nm. The temperature at the end of rough rolling is lowered so that the center of the steel product/transfer bar is at about 1040° C., the temperature of no recrystallization corresponding to a Zener pinning pressure of 0.1 MPa obtained by TiN—NbC composites of 32 nm mean size with a mean interparticle spacing of 200-300 nm. Electron energy loss spectroscopy has confirmed growth of NbC on pre-existing TiN. The limiting austenite grain size by TiN—NbC composite precipitates is less than 30 microns, which requires less pancaking in finish rolling, resulting in thicker gage (>16 mm).

FIG. 9 is a montage that relates interparticle spacing of nano-scale TiN—NbC composites to titanium and nitrogen content in the base chemical composition, which is mapped on the equilibrium solubility product for TiN precipitation as a function of temperature. The montage represents a comprehensive data base on interparticle spacing of TiN obtained in line pipe steel in which nitrogen content is varied. The interparticle spacing of TiN is in the 200-250 nm range when nitrogen content is lowered to 40 ppm, titanium is added in the stoichiometric requirement to combine with all the nitrogen. High number density of TiN—NbC composite precipitates occurs in the size range of 25-35 nm. By contrast, when nitrogen content is raised to 75 ppm, the interparticle size is large at about 550 nm, and the particle size is coarse (80 nm size).

FIG. 10 is an optical micrograph showing austenite grain size in the transfer bar of Steel-D quenched after shearing. The austenite grain size is about 48-55 microns. This is in agreement with Zener limiting austenite grain size, based on measured values of precipitate size and interparticle spacing of Steel-D, shown in FIG. 9.

FIG. 11 is Kozazu's diagram, inter-relating rolling reduction and austenite grain size with surface to volume ratio, Sv factor, of pancaked austenite grain size, to which the final structure and properties can be related. Kozazu's diagram shows that a large rolling reduction (70 percent) is required to pancake coarser austenite grain of 70 micron compared with lower rolling reduction (<50 percent) required to pancake finer austenite grain of 30 micron grain size to achieve the same surface to volume ratio, i.e., Sv factor.

FIG. 12 is a photograph compiled from Electron Energy Loss Spectroscopy (EELS) data of nano-scale TiN—NbC precipitates observed in nano-scale precipitation engineered high grade line pipe steel processed by conventional hot strip rolling. The epitaxial growth of NbC on faces of the TiN cubic precipitates can be clearly seen in Steel-C(X-90 grade).

FIG. 13 shows elemental mapping from EELS data of the TiN—NbC composite precipitates, shown in FIG. 12. These results show unambiguously epitaxial growth of NbC on pre-existing TiN.

FIG. 14 illustrates the application of nano-scale precipitation engineering of TiN—NbC composite precipitates for austenite grain size control in near net shape processing for a typical lay out of mill design with three roughing stands.

FIGS. 15A, B and C show metallographic evidence of grain coarsening of fine austenite grains obtained by traditional static recrystallization after 25% deformation followed by isothermal holding for 3, 60 and 240 seconds, respectively, at 1000° C.

FIG. 16 is a plot showing the predicted grain coarsening behavior of grain refined austenite grains from the end of rough rolling to start of finish rolling with respect to time with for a steel product with 0.03 wt % solute niobium for the temperature window and the specific applicable cooling rate.

FIG. 17 depicts the evolution of transformation texture of recrystallized austenite grains showing cubic texture in comparison with the pancaked austenite grains. The undesirable rotated cube texture from recrystallized cubic austenite can be suppressed by deforming the austenite grain to brass or copper texture before transformation.

FIG. 18 depicts a process sheet in accordance with the processes of embodiments of the present invention for producing thicker gage steel product in a conventional hot strip mill.

FIG. 19A depicts a flow chart depicting the prior art process and a transmission electron microscopy (TEM) image of delayed strain induced precipitates of NbC that result therefrom.

FIG. 19B depicts a flow chart depicting a process in accordance with embodiments of the present invention and a TEM image of growth of NbC on pre-existing TiN to form TiN—NbC that suppresses delayed strain induced precipitation of NbC.

FIG. 20 is a schematic diagram that illustrates the effect of nitrogen content on interparticle spacing of TiN—NbC composite and delayed strain induced precipitation of NbC.

FIG. 21 depicts a 3D atom probe tomography reconstruction Nb atomic map and a table with the corresponding average chemical composition of the conventional hot strip rolled X90 example of Steel-C in accordance with the present invention. The solute niobium retained in the matrix is 0.031 wt %.

FIG. 22A depicts a TEM image of fine acicular ferrite obtained in X-90 strip showing an acicular ferrite grain size of 1-3 μm, produced in accordance with embodiments of the present invention.

FIG. 22B depicts a TEM image of fine acicular ferrite of about of 1 μm obtained in X-90 strip along with a fine dispersion of fine Martensite-Austenite, e.g. MA, product of <0.5 μm, produced in accordance with embodiments of the present invention.

FIG. 23 shows band contrast map from Electron Back Scatter diffraction (EBSD) of X-90 strip of 16.4 mm thickness, depicting the target microstructure in accordance with embodiments of the present invention, showing pancaked austenite thickness of under 10 microns, about 6 microns near the surface region and about 8.5 microns in the center region of the strip.

FIG. 24 shows EBSD characterization of the center of X-90 strip processed by nano-scale precipitate engineering of TiN—NbC composite in accordance with embodiments of the present invention. The EBSD image on the left shows morphological microstructure of acicular ferrite grains of 1-3 microns in pancaked austenite grain of thickness of about 8.5 microns. The EBSD image on the right shows high density and uniform dispersion of high angle boundaries above 45 degrees misorientation that gives crystallographic fine domain size which suppresses brittle fracture. The EBSD image on the right also shows sparse dispersion of rotated cube texture as delineated by the regions shaded darker than the others. This EBSD image gives the target morphological and crystallographic structure to obtain high strength and good fracture toughness at low temperature.

FIG. 25 depicts a comparison of chemistry, processing, precipitates and {100}<011> cleavage texture, DWTT, and Crack Tip Opening Displacement (CTOD) test between Trial-1 (traditional strain induced precipitation) and Trial-2 (austenite grain refining by Nano-scale precipitate engineering of TiN—NbC in accordance with embodiments of the present invention). The images in part III depict EBSD results for the 32 mm gage plate rolling of Trial-1 (image on the left) exhibiting rotated cube texture associated with conventional processing based on strain induced precipitation of NbC in comparison with Trial-2 (image on the right) based on nano-scale precipitate engineering of TiN—NbC composite, in accordance with embodiments of the present invention, that suppresses the formation of undesirable rotated cube texture. The shaded regions in the picture on the left correspond to a high intensity of undesirable rotated cube texture from Trial-1 (processed according to prior art), while the image on the right shows suppressed rotated cube texture from Trial-2 processed according to the embodiments of the present invention.

FIG. 26 depicts a plot of percentage shear area from a DWTT (as measured in accordance with API RP 5L3 (Apr. 1, 1996)) plotted as a function of test temperature for 32 mm K-60 plate processed by conventional strain induced nucleation of NbC (Trial-1), and for K-60 plate and X-90 strip processed by nano-scale precipitate engineering of TiN—NbC composite methods in accordance with embodiments of the present invention (Trial-2).

FIG. 27 depicts a comparison between traditional processes (left) and embodiments of processes in accordance with embodiments of the present invention (right).

FIG. 28 is a schematic depiction of the functional role of nano-scale precipitate engineering of TiN—NbC composite in austenite processing.

FIG. 29 is a schematic depiction of the functional role of solute niobium in promoting crystallographic high angle boundaries through its hardenability effect to suppress brittle fracture.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Embodiments of the invention are based on strain induced growth of NbC on pre-existing TiN with short interparticle spacing of about 200-300 nm to form TiN—NbC composites of adequate size of about 30 nm at negligible undercooling (<5° C.) so that the temperature of no recrystallization of austenite grains (Tnr) is promoted at about 100° C. higher than in conventional processing of niobium microalloyed steels in which temperature of no recrystallization occurs by strain induced nucleation of NbC precipitates at large undercoolings (>100° C.). By raising temperature of no recrystallization, the undesirable rotated cube texture that promotes brittle fracture is suppressed by pancaking the austenite grains before transformation. Delayed strain induced precipitation renders it difficult to control the transformation temperature in downstream processing and the resulting structure and properties. By controlling the interparticle spacing of TiN—NbC, the undesirable delayed strain induced precipitation of NbC in austenite is suppressed.

FIGS. 1A and 1B are electron backscatter diffraction images of 13% Cr-5% Ni-2% Mo super-martensitic stainless steels without and with 0.1 wt percent niobium addition. These steel specimens were identically processed, solution treated at 1050° C. and air-cooled. The detailed chemical compositions of the two steels are given in Table 2.

TABLE 2 Chemical composition of two 13%Cr—5%Ni—2%Mo steels without and with niobium in wt percent Steel grade C Si Mn P S Cr Ni Mo N Nb Ti 13Cr5Ni2Mo 0.020 0.42 0.51 0.016 0.004 12.59 5.01 1.90 0.013 — 0.0062 13Cr5Ni2MoNb 0.022 0.41 0.48 0.016 0.006 12.91 5.16 2.05 0.010 0.11 0.0043

The steel without niobium addition but with titanium exhibits an austenite grain size of 80 microns, which shows that titanium addition alone is not effective in refining austenite grain size. But with the addition of 0.1 wt percent niobium, the austenite grain size is significantly decreased to 35 microns. The white lines in FIG. 1 delineate the austenite grain boundaries.

FIG. 2 shows TEM images of composite precipitates of TiN—NbC observed in the steel with a niobium addition. The mean interparticle spacing of composite precipitates is 231 nm and the mean particle size is increased from 15 nm for TiN to 30 nm for the TiN—NbC composite particles. The pinning pressure exerted by the particles of TiN—NbC is 0.08 MPa, which counteracts the driving force for grain coarsening of austenite of grain size 35 microns as shown in Table 3.

TABLE 3 Calculated limiting austenite grain size by TiN and TiN—NbC in 13Cr5Ni2Mo steel with 0.1 wt. percent niobium Zener Limiting Interparticle pinning austenite Size Spacing Volume pressures grain size Precipitate (nm) (nm) fraction (MPa) (μm) TiN 15 231 0.00015 0.02088 134 TiN—NbC 30 231 0.00123 0.08028 35

The area occupied by particles on the boundary must be recreated before the boundary moves, and this is the energy preventing grain coarsening and is referred to as Zener drag. Solute atoms piling up at the interface exert a drag force on boundary mobility, which is referred to as solute drag. Solute drag of niobium is less pronounced in the high temperature window. Zener drag is increased as the interparticle distance is reduced and the particle size becomes bigger for a given volume fraction of precipitate, which is determined by the thermodynamic potential for precipitation for a given steel composition. TiN and TiN—NbC precipitates occur on a nano-scale and therefore engineering the size and dispersion of nano-scale precipitates is termed “engineering nano-scale precipitates for pinning grain boundaries.”

FIG. 3 shows that TiN particles by themselves are not effective in pinning austenite grains. The epitaxial growth of NbC on pre-existing TiN effectively increases the particle size from 15 to 30 nm, with the corresponding increase in Zener pinning pressure nearly threefold (2.5 times) compared with TiN alone and thus decreases the limiting austenite grain size. This composite precipitate involves growth of NbC on pre-existing TiN, which is to be distinguished from MX type precipitates reported in the literature and previous U.S. Pat. No. 6,899,773.

FIG. 4 and FIG. 5 provide a summary of the prior work on titanium-niobium microalloying, which shows that if nitrogen is controlled under 40 ppm and titanium additions are made to the stoichiometric requirement of N to form TiN, a high number density of TiN can be promoted. On reaching the temperature where thermodynamic potential for precipitation of NbC occurs, NbC will start to grow on pre-existing TiN. This work was reported by S. V. Subramanian, F. Boratto, J. J. Jonas and C. M. Sellars and published in the Proceedings of International Symposium on “Microalloyed Bar and Forging Steels” edited by Mike Finn, CIM, held at Hamilton, Ontario, Canada, Aug. 26-29, 1990, pp. 120-136.

FIG. 4 shows a TEM image of TiN—NbC composite precipitates. The energy dispersive analysis and EELS (Electron energy loss spectrum) have confirmed that NbC precipitates grow epitaxially on the faces of cuboidal precipitates of TiN (with the NaCl crystal structure.)

FIG. 5 shows that growth of NbC on pre-existing TiN can be detected by an increase in flow stress during hot torsion simulation of rolling. The temperature corresponding to the increase in flow stress detected in the hot torsion experiment is found to coincide with the equilibrium temperature of precipitation of NbC, as shown in Table 4. The implication is that the volume fraction of NbC growing on pre-existing TiN can be determined by the thermodynamic potential for precipitation of NbC.

TABLE 4 Effect of chemical composition on precipitation kinetics of Ti—Nb microalloyed steel during hot torsion simulation of rolling; Growth of NbC on pre-existing TiN occurs close to the equilibrium temperature for precipitation of NbC, which can be detected by flow stress increase during hot torsion. T_(NbC) Sellars Model Base Chemistry Equilibrium RLT RST T_(nrHot Torsion) Steel C N Ti Nb ° C. ° C. ° C. ° C. Fe—Nb—C 0.15 0.005 0 0.031 1130 955 922 960 Line pipe A 0.029 0.0023 0.013 0.055 1010 898 880 1010 Line pipe B 0.026 0.0027 0.020 0.057 1020 900 885 1016 Line pipe C 0.016 0.0018 0.013 0.049 960 850 836 994 Line pipe G 0.027 0.0035 0.014 0.099 1060 936 923 1060 Line pipe K 0.027 0.0019 0.014 0.093 1050 936 923 1040 RLT = Recrystallization Limit Temperature RST = Recrystallization Stop Temperature T_(nr) = Temperature of No Recrystallization

The breakthrough in austenite grain size control upstream arose out of observations in super-martensitic stainless steel, where a titanium addition by itself did not produce a fine austenite grain size. But when combined with niobium additions, a fine austenite grain size was obtained. This was caused by TiN—NbC composite precipitates, formed by NbC growing on pre-existing TiN. The number density of precipitates was controlled by the TiN. By increasing the size of the precipitates by growth of NbC on pre-existing TiN, it is possible to increase the pinning pressure of precipitates to arrest austenite grain boundary movement at the required austenite grain size. This discovery underlies the present invention of nano-scale precipitate engineering. The austenite grain size is controlled by interparticle spacing of TiN precipitates and the size of the precipitates, each of which can be independently controlled by design of the base steel composition. A high number density of TiN is promoted when the precipitation occurs in the matrix at low temperature, which calls for lowering the nitrogen content and adding titanium to the stoichiometric requirement to form TiN, providing one atom of titanium for every one atom of nitrogen. By lowering nitrogen to less than 40 ppm and adding titanium to the stoichiometric requirement of about 0.014 wt percent, the interparticle spacing was found to be around 200 nm and the TiN precipitate size was found to be in the 10-15 nm range. The Zener pinning pressure on the boundary is relatively small, capable of arresting austenite grains of about 80 microns from coarsening. The pinning pressure can be increased by growing NbC precipitates on TiN, thereby increasing the size of the composite precipitates of TiN—NbC. This requires lowering the temperature window of roughing so that the thermodynamic potential for growth of NbC on pre-existing TiN is obtained to form TiN—NbC composites. By increasing the particle size while retaining the same interparticle spacing, the pinning pressure is increased to arrest the finer austenite grains from coarsening. By increasing the mean size of precipitates to about 30 nm with an interparticle spacing of about 220 nm, the Zener pinning pressure is increased to prevent austenite grain size of 30 microns from coarsening.

The concept of nano-scale precipitation engineering to arrest grain coarsening is illustrated in the schematic diagram given in FIG. 6. The upstream processing of austenite for austenite grain size control requires a high number density of precipitates with short interparticle spacing and adequate precipitate size with good dispersion to apply adequate pinning pressure to prevent coarsening of fine grains of austenite obtained at the end of roughing. This innovation relates to product-process integration, where refinement of austenite grain size in upstream processing by Zener pinning by TiN—NbC composite precipitates of grain refined austenite to prevent grain coarsening is used to reduce total rolling reduction in finish rolling downstream to produce thicker gage product.

Nitrogen is controlled to promote the formation of TiN precipitates at lower temperatures. The resulting finely-dispersed nano-precipitates of TiN then act as scaffolds for the epitaxial formation of NbC, thereby raising the volume fraction of dispersed composite precipitates by a factor of about 3×. This is sufficient to hold the austenite grain size to about 30 micrometer size in low nitrogen steel compared to 60 microns for higher nitrogen steel. The advantage here lies in the reduced austenite grain size, permitting the application of a reduced rolling reduction during final processing and the consequent ability to produce thicker gages of higher strength material (X-70, X-80, X-90, X-100) compared with high nitrogen steel, which requires heavy rolling reduction that limits final gage of the product.

FIG. 7 is the flow diagram of product-process integration of the prior art technology without any intentional control of austenite grain size upstream and its consequence on heavy pancaking downstream resulting in thin gage product. A high nitrogen content in the base composition results in coarse precipitates of TiN with a large interparticle spacing of 550 nm. TEM characterization shows the coarse precipitate of TiN. Rough rolling is carried out in a temperature window above the equilibrium temperature for precipitation of NbC. Thus, Zener limiting austenite grain size is 62 microns, as shown in the Table in FIG. 7. Therefore, heavy pancaking is required to obtain consistently good DWTT performance.

By comparison, FIG. 8 is the flow diagram of product-process integration of the present invention based on austenite grain size control by engineering size and spacing of TiN—NbC composite precipitates and its consequence on reduced rolling reduction downstream, resulting in production of thicker gage product. A low nitrogen content of 40 ppm with stoichiometric addition of Ti to combine with all nitrogen promotes in high number density of TiN precipitates with an interparticle spacing of 220 nm. The temperature window of roughing is lowered below the equilibrium temperature for precipitation of NbC to promote growth of NbC on pre-existing TiN, which is confirmed by EELS characterization of TiN—NbC composite precipitates. The limiting austenite grain size is below 30 microns. Therefore, less rolling reduction is applied to produce thicker gage product.

The technology of nano-scale precipitation engineering of TiN—NbC composites involves two microstructural parameters. The first is the interparticle spacing. The second is the particle size. This invention is based on the discovery that TiN—NbC composites offer a window of opportunity to control interparticle spacing through optimum TiN distribution and the size of the particle by epitaxial growth of NbC on pre-existing TiN particles. The first step is to engineer a high number density and uniform dispersion of TiN particles. This is done by promoting nucleation of TiN in austenite at lower temperatures through control of the base steel chemical composition. Since the precipitates occur on a nano-scale, it is essential to characterize the precipitates by transmission electron microscope. The well-known carbon replica technique is used in this work to extract the precipitates occurring in benchmarked steels. FIG. 9 shows a comprehensive database of four bench marked steels in which nitrogen content is varied under different mill processing conditions. The chemical compositions of the four steels are given in Table 5.

TABLE 5 Effect of varying nitrogen content on thermodynamic potential for precipitation of TiN, and its consequence on ppt size and Zener pinning pressure, and Zener limiting austenite grain size. The interparticle spacing of TiN and the size of TiN—NbC are measured values. TiN Inter- TiN—NbC Volume fraction N Ti C Nb particle Ppt size of TiN—NbC Zener pinning Limiting austenite I.D. (wt %) (wt %) (wt %) (wt %) spacing in (nm) (nm) at 1000° C. Pressure (MPa) grain size (μm) A 0.0075 0.015 0.06 0.09 553 83 0.00176 0.044 62 B 0.0035 0.014 0.07 0.08 218 32 0.0016 0.108 26 C 0.0040 0.015 0.05 0.09 221 32 0.00159 0.104 27 D 0.0055 0.012 0.048 0.067 397 52 0.00117 0.047 59

Steel-A with the highest nitrogen content of 0.0075 wt. percent exhibits a large mean interparticle spacing of about 550 nm compared with Steels-B and C with a low N content of 35-40 ppm, which exhibit a mean interparticle spacing of about 220 nm. Steel-D with intermediate nitrogen content of 55 ppm exhibits an intermediate interparticle spacing of about 400 nm. Clearly, the interparticle spacing of 220 nm can be achieved by lowering nitrogen content to or below 40 ppm and adding titanium to the stoichiometric requirement to tie up all the nitrogen. The precipitate size of TiN—NbC of the highest nitrogen Steel-A is 83 nm, which gives Zener limiting austenite size of 62 microns. By comparison, Steel-B and Steel-C with low nitrogen give Zener limiting austenite grain size of about 27 microns. FIG. 10 shows the austenite grain size measured in the center of a thick transfer bar of 53 mm of Steel-D, quenched after rough rolling with an intermediate nitrogen content of 55 ppm and an interparticle spacing of 397 nm. The predicted Zener limiting austenite grain size is 59 microns, which compares well with the measured value of 55 microns, which validates the approach. Thus, nano-scale precipitation engineering offers a sound metallurgical basis for controlling austenite grain size during upstream processing of austenite.

In conventional processing of conventional nitrogen-bearing niobium microalloyed steel (0.005-0.008 wt. percent nitrogen), roughing is carried out where there is no thermodynamic potential for precipitation of NbC. The loss of niobium by excessive growth of NbC on pre-existing TiN particles is reduced by minimizing the time of processing in the mill. In the case of nano-scale precipitation engineering of TiN—NbC composites with finer interparticle spacing, it is even more critical to prevent depletion of solute niobium in the matrix by accelerated cooling upstream between the end of roughing and the start of finish rolling. It is essential to control the finish rolling entry temperature below the temperature of no recrystallization in order to avoid rolling in the partial recrystallization regime, which requires accelerated cooling. Thus, accelerated cooling is required to prevent depletion of solute niobium by precipitate growth, subsequent to pinning the austenite grains of the required size in higher grade line pipe steel.

TABLE 6 Effect of austenite grain size (GS) and percent reduction below temperature of no recrystallization (T_(NR)) on Sv factor and ferrite grain size. Ferrite GS Sv Factor Austenite GS % Reduction (um) mm²/mm³ (um) below T_(NR) 9 80 40 60 9 80 30 30 9.4 70 55 60 9.4 70 35 30 11 60 70 60 11 60 40 30

Table 6 is extracted from Kozazu's diagram in FIG. 11, which illustrates the benefit of austenite grain refinement before pancaking in reducing the rolling reduction below the temperature of no recrystallization to achieve the same surface to volume ratio. Thus, by reducing the austenite grain size from 40 to 30 microns, the rolling reduction can be decreased from 60 to 30 percent to attain the same Sv factor of 80 mm²/mm³ in order to obtain ferrite grain size of 9 micrometers and consequently the gage (thickness of final product) can be significantly increased. It is well established that by refining the austenite grain size upstream, excellent strength and fracture properties can be obtained in thicker gage product. Wenjin Nie et al. have demonstrated the importance of austenite grain size control on final DWTT properties of heavy thick X-80 pipe line steels (Advanced Materials Research Vols. 194-196, 2011, pp. 1183-1191).

Suppression of Delayed Strain Induced Precipitation:

As shown in FIG. 18, a process according to another embodiment of the invention includes two stages in roughing. Stage-1 of roughing is the same as in prior art to refine austenite grains below 30 microns. Stage-1 can include multiple passes, see e.g., FIG. 18 R1-R7, when the surface temperature is in a window between 1080° C.-1150° C. During Stage-1, the temperature gradient between the surface and the core is less, therefore if the surface temperature in Stage-1, pass R1 is about 1150° C. the core temperature will be slightly higher. As Stage-1 continues into the later passes, the temperature gradient between surface and core will increase. Stage-2 of roughing involves holding for about 100-120 seconds and cooling for temperature homogenization in roughing, when grains will coarsen during roughing. Then Stage-2 includes applying at least one heavy pass reduction (when the temperature in the core is about 1040° C.), e.g. schematically shown as R8, austenite grains are refined again and pinned by strain induced growth of NbC on pre-existing TiN with short interparticle spacing of about 200 nm to exert a pinning pressure of 0.1 MPa to suppress recrystallization and grain coarsening in the center of the transfer bar, thereby preventing undesirable rotated cube texture formation from occurring in the center while ensuring large Sv factor due to austenite grain size control before pancaking. It is contemplated that more than one heavy pass can be used in the second stage of roughing (e.g. an R9 pass) as in heavy gage (32 mm) plate rolling. The process of FIG. 18 can be performed using a conventional hot strip mill. For example, the conventional hot strip mill includes one or two rolling stands and the rolling direction for the rolling stands is reversible, which enables multiple passes in single rolling stand by reversible directional rolling.

With reference now to FIG. 19A, TEM images of NbC precipitates in conventional niobium microalloyed steel processed by Steckel Mill rolling using strain induced nucleation of NbC to control Tnr are shown. In FIG. 19B, TEM images of NbC precipitates occurring in coil processed by nano-scale precipitate engineering of TiN—NbC composite in accordance with the embodiments of the present invention are shown. As shown by the comparison between FIGS. 19A and 19B, in the absence of nitrogen control, the occurrence of TiN with large interparticle spacing promotes delayed strain induced precipitation of NbC. The interparticle spacing of delayed strain induced precipitates of NbC is difficult to control. The spacing of precipitates range from 20-40 nm and their size extend from 5-15 nm. The occurrence of high number density of NbC precipitates with short inter-spacings renders it difficult to control residual niobium in the matrix and consequently it is difficult to suppress polygonal ferrite. According to the prior art by Bal et al., the microstructure exhibits 30% polygonal ferrite with 70% acicular ferrite. By comparison, in embodiments of the present disclosure, the control of interparticle spacing of TiN—NbC composite is controlled by TiN in the range of 200-300 nm, which suppresses delayed strain induced precipitation of NbC and promotes 100% acicular ferrite morphology.

As shown in FIG. 20, a schematic diagram illustrates the importance of control of interparticle spacing upstream on the evolution of delayed strain induced precipitation of NbC. The low-nitrogen control of the present invention is shown in the top diagram, and the high nitrogen in traditional processes is shown in the bottom diagram. The experimental data on delayed strain induced precipitation of NbC is analyzed in detail in the paper entitled “Suppression of strain induced precipitation of NbC by epitaxial growth of NbC on pre-existing TiN in Nb—Ti microalloyed steel” by Xiaoping Ma, Chengliang Miao, Brian Langelier and Sundaresa Subramanian, Materials and Design, 132, (2017), pages 244-249. By controlling the interparticle spacing of TiN—NbC composite through TiN, it is found that delayed strain induced precipitation of NbC can be totally suppressed, thereby ensuring adequate solute niobium in the matrix for control of target microstructure of acicular ferrite produced at low temperature of transformation. The control of residual solute niobium is confirmed by atom probe studies to be about 0.03 wt % or more, which promotes acicular ferrite of 1-3 microns by shear transformation, see FIG. 21. The structure exhibits high density and dispersion of high angle boundaries that suppress brittle fracture.

FIGS. 22A-22B depict TEM images of acicular ferrite morphology obtained in X-90 strip processed by nano-scale precipitate engineering of TiN—NbC composite. FIG. 22A depicts a TEM image of fine acicular ferrite obtained in X-90 strip showing acicular ferrite of 1-3 micron size. FIG. 22B depicts a TEM image of fine acicular ferrite obtained in X-90 strip showing fine Martensite-Austenite (MA) product of <0.5 μm size, which is well below the critical size to initiate brittle fracture.

FIG. 23 shows pancaked austenite thickness in the region close to the surface and in the center of the strip. The pancaked austenite thickness near the surface is about 6 microns, and that in the center is 8.5 microns. The acicular ferrite grain size within the austenite grain is 1-3 microns. The uniformity of microstructure confers consistent strength and toughness properties.

FIG. 24 depicts the target microstructure in accordance with embodiments of the present invention, showing acicular ferrite of 1-3 microns in pancaked austenite of thickness of about 8.5 microns. The left hand micrograph of FIG. 24 includes measurements showing 8.6 microns and 8.4 microns. There is no polygonal ferrite unlike the results reported in prior art by Bal et al. (20% polygonal ferrite). The analysis of crystallographic data of the same sample summarized in the right hand image shows a high density of high angle boundaries, with misorientation above 45 degrees that arrest micro-cracks before they grow to critical Griffith crack length to initiate brittle fracture. The intensity of rotated cube texture is suppressed as revealed by sparse distribution of shaded area corresponding to rotated cube texture. These are significant advantages of embodiments of the present invention.

Control of Rotated cube texture:

In traditional microalloyed steel, the center of heavy gage plate as well as strips of thermo-mechanically rolled steels is found to exhibit rotated cube texture, which promotes cleavage brittle fracture. This is caused by transformation of cubic texture associated with recrystallized austenite grains. By work-hardening the austenite before transformation, it is possible to suppress the occurrence of undesirable rotated cube texture. In order to work harden the austenite in the center of thick section, it is essential to suppress recrystallization by raising temperature of no recrystallization. In the prior art, the temperature of no recrystallization by strain induced nucleation of NbC occurs at large undercoolings, and hence, Tnr is typically lower than 960° C. Since the center of thick section is about 100° C. higher than the surface, it is difficult to suppress rotated cube texture from occurring in the center of thicker gage unless Tnr is raised substantially. Nano-scale precipitate engineering of TiN—NbC composite, in accordance with embodiments of the present invention, raises Tnr by about 100° C. by strain induced growth of NbC on pre-existing TiN at negligible undercoolings, see Table 4. Since nano-scale precipitate engineering of TiN—NbC raises temperature of no recrystallization significantly (by more than 100° C.), it can be used to advantage to arrest recrystallization of austenite occurring in the center of thick section which is about 100° C. higher than the surface temperature, thereby work-harden the austenite to modify the cubic structure of austenite to brass or copper before transformation in order to eliminate rotated cube texture.

EXAMPLES

The principal differences in the processing of higher niobium steels between the prior technology without austenite grain size control upstream and the technology of the present invention based on austenite grain size control are examined in further detail and their consequence on product in terms of gage thickness and properties are highlighted in the following examples.

Example 1: Plate Rolling

Steel-A with High Nitrogen Content

Steel-A is representative of prior technological practice, where a higher nitrogen content of 75 ppm gives coarse TiN particles with large interparticle spacing of 550 nm. Roughing is carried out in the temperature window where there is no thermodynamic potential for precipitation of niobium carbide. Thus, the austenite grain size entering finish rolling is 60-80 microns. This then requires heavy pancaking below the temperature of no recrystallization. Thus, the final gage is generally limited to 16 mm. Typical property results obtained from 10 mm gage are reproduced below in Table 7:

TABLE 7 End of Limiting austenite N Ti C Nb roughing grain size by TiN 0.0075 0.015 0.06 0.088 1100° C. 90 microns DWTT % SA at −7° C.: 100%; CVN toughness at −7° C.: 140 Joules; Yield Strength/Rp0.5: 610 MPa; Ultimate Tensile Strength/Rm: 714 MPa

Example 2

Steel-E: Plate Rolling with Low Nitrogen Content:

Steel-E has a lower nitrogen content (40 ppm) with titanium and niobium addition comparable to Steel-A. The low nitrogen and stoichiometric addition of titanium to combine with nitrogen to form TiN has produced a high number density of TiN with a mean interparticle spacing of 220 nm. This steel was processed under two distinctly different conditions. The first set of conditions was where the rough rolling window was similar to Steel-A, that is where there is no thermodynamic potential for NbC precipitation to occur. Under these conditions, TiN particles alone are not able to develop pinning pressure adequate to pin a fine austenite grain size. Thus, the resulting coarse austenite grain size warrants heavy rolling reduction, which is not possible to achieve in 22 mm gage thickness. As a consequence, the final product fails as percentage shear area in the DWTT specimen is lowered to 55 percent at −15° C. (See Table 8).

Steel-E was also processed under a second set of conditions, more than one heavy pass of about 25% rolling reduction was carried out below the dissolution temperature of NbC. In this case, NbC grows on pre-existing TiN to increase the particle size so that the pinning pressure is increased to prevent austenite grain coarsening above 30 microns.

Once austenite grain size is refined at the entry to finish rolling, less rolling reduction is required in finish rolling in accordance with Kozazu's diagram in FIG. 11 to obtain adequate surface to volume ratio to obtain fine grains in the final product. In this case, 100 percent shear area on the fracture surface in DWTT is obtained. This example shows that TiN by itself cannot grain refine austenite even though the interparticle spacing may be fine unless the particle size is increased by growth of NbC on pre-existing TiN particle. This example demonstrates the importance of lowering the temperature window of roughing to promote growth of NbC on pre-existing TiN to limit austenite grain coarsening at the end of roughing in order to produce 22 mm gage with excellent DWTT performance.

Steel-E

22 mm thick gage—X80 Plate: (nitrogen 0.004, titanium 0.016, carbon 0.05, niobium 0.1)

Effect of processing temperature window on low nitrogen and high niobium steel.

Effect of rough rolling in the temperature window with and without thermodynamic potential for precipitation of NbC.

TABLE 8 #1 Condition #2 Condition (Roughing without (Roughing to promote Heat NbC growth on TiN) NbC growth on TiN) A_(k) −20° C./Joules 328 372 DWTT (−15° C) 55 98 average SA %

Example 3 (Conventional Hot Strip Rolling):

Steel-D with Intermediate Nitrogen Content

Steel-D presents a case, where nitrogen content is at an intermediate level of about 55 ppm and therefore the mean interparticle distance is 390 nm. Though the temperature of finish rolling promoted growth of NbC on pre-existing TiN, the TiN—NbC composite did not have adequate pinning pressure to arrest austenite grains finer than 59 microns, see Table 9. This is partly due to low niobium content, i.e. 0.067 wt. percent. Thus, the strip rolled to 20 mm gage thickness exhibited 100 percent shear only at −10° C. and above, see Table 10a and 10b.

Steel-D: 20 mm thick gage X80 Strip

High nitrogen and lower niobium with rough rolling in the temperature regime where there is thermodynamic potential for precipitation of NbC.

TABLE 9 Limiting Pancaking End of austenite Pancaking austenite N Ti C Nb roughing grain size reduction grain 0.0055 0.012 0.048 0.067 980° C. 59 microns 62.6 20 microns

TABLE 10a Yield Ultimate tensile Yield Total strength/Rp0.2 Strength/Rm ratio elongation/% 588 MPa 670 MPa 0.88 48

TABLE 10b Test DWTT Shear Area % Temperature/ Charpy V-notched T L 45° ° C. toughness/Joules direction direction direction 0 — 95 95 95 −10 — — 100 90 −20 485 70 95 100

Example 4 Conventional Hot Strip Rolling

In-depth characterization of Steel-C has confirmed that the interparticle spacing of TiN is 220 nm. Steel-C represents low nitrogen content, with optimized addition of titanium to promote high number density and uniform dispersion of TiN with an interparticle spacing of 220 nm. The end of roughing is in the temperature window where thermodynamic potential for precipitation of NbC occurs.

The process steps involve two stages of roughing. The first stage of roughing is to refine the austenite grain size by static recrystallization in the first few passes of rough rolling in the high temperature window where there is no significant thermodynamic potential for precipitation of NbC. The second stage of roughing is at a much lower temperature, e.g. 1040° C. in the core and a corresponding surface temperature of 940° C. at the surface in order to obtain significant thermodynamic potential for strain induced growth of NbC on pre-existing TiN. This requires holding for about 100 seconds after the first stage of roughing for temperature homogenization through the section thickness. In this context, temperature homogenization refers to the temperature gradient across the section thickness decreasing. The target temperature is to promote Tnr in the center of the section, which is about 1040° C., the corresponding surface temperature will be considerably less (940° C.). Thus, a heavy rolling pass reduction (>25%) is applied as the last pass of the second stage of roughing in order to promote Tnr in the center of transfer bar by strain induced growth of NbC on pre-existing TiN to promote adequate pinning pressure due to TiN—NbC precipitates of about 0.1 MPa, on austenite grain boundaries to arrest recrystallization and prevent grain coarsening during subsequent cooling to enter finish rolling with a center temperature at about 920° C. Without the rolling reduction of the second stage of roughing at a low temperature to promote strain induced growth of NbC, austenite grains will coarsen from less than 30 microns to 60 microns and above. By Zener pinning of austenite grains, grain coarsening of austenite beyond 30 microns is prevented. Upon pancaking austenite grains with a large rolling reduction of 70%, the thickness of pancaked austenite grains is limited to 8 microns and below, which gives excellent toughness at low temperature.

TEM-EELS characterization of TiN—NbC precipitates shown in FIGS. 12 and 13 confirms epitaxial growth of NbC on pre-existing TiN particles. This steel exhibits remarkable toughness at very low temperature (−40° C.), see Tables 11a and 11b. The steel exhibits uniformity of microstructure which is less prone to anisotropic properties due to unfavorable austenite grain texture development.

Steel-C: 16.4 mm thick gage X90 Strip.

Low nitrogen and higher niobium with rough rolling in the temperature regime where there is thermodynamic potential for precipitation of NbC.

TABLE 11a Yield Ultimate tensile Yield Total Uniform strength/Rp0.2 Strength/Rm ratio elongation/% elongation/% 670 MPa 800 MPa 0.84 17 5.6

TABLE 11b Testing DWTT Temperature/ Charpy V-notched (T direction) ° C. toughness/Joules Shear area % 10 313 0 302 100 −10 300 100 −20 315 100 −40 318 100 −60 329

Example 5: Compact Strip Processing and Thin Slab Processing

There are different mill designs available for compact strip processing. Nano-scale precipitate engineering of TiN—NbC composites offers a generic platform for preventing austenite grain coarsening by controlling interparticle distance by TiN, and particle size by NbC growing on the pre-existing TiN. In near net shape processing, in some cases, the transfer bar is reheated for the purpose of temperature homogenization, when the austenite grains inevitably coarsen in the absence of second phase particles. The technology of nano-scale precipitation engineering offers a sound basis for pinning austenite grain boundary with TiN—NbC composite precipitates at the end of roughing, and also during reheating. This process can be combined with accelerated cooling to prevent depletion of solute niobium by excessive growth of NbC, over and above the composite particle size required to prevent grain coarsening of austenite of a specific grain size. Trials of nano-scale precipitation engineering in a mill with two roughing stands and accelerated cooling at 4° C./s have given uniformity of microstructure, which is beneficial in achieving consistent strength and fracture properties. Delayed strain induced precipitation of NbC is reported to cause unpredictable mill loading in near net shape processing. The control of interparticle spacing of TiN—NbC offers a novel metallurgical solution to overcome this problem by suppressing delayed strain induced precipitation of NbC.

Example-6: Application of Nano-Scale Precipitate Engineering of TiN—NbC Composite in Plate Rolling of 32 mm K-60 E2 Grade

Experimental results are presented that confirm the benefits of nano-scale precipitate engineering of TiN—NbC composite by strain induced growth of NbC on pre-existing TiN over conventional processing by strain induced nucleation of NbC in processing 32 mm gage plate, using low nitrogen base chemistry in Example-6.

Trial-1 and 2 involved two alternative methods of processing respectively but using low nitrogen base chemistry. Steel chemistry of Trial-1 contained 0.04C, 0.0022N, 0.087Nb and 0.017Ti and that of Trial-2 contained 0.035C, 0.0029N, 0.083Nb and 0.016 Ti. Each trial was conducted with low nitrogen (<30 ppm), Ti (0.015 wt %) at the stoichiometric requirement to tie up all N, low carbon (0.04%) and high niobium (0.085 wt %). Plate rolling in Trial-1 was based on promoting strain induced nucleation of NbC, whereas rolling in Trial-2 was based on promoting strain induced growth of NbC on pre-existing TiN. FIG. 25 summarizes TEM and EBSD results of Trial-1 and Trial-2.

A comparison of the DWTT of Trial-1 (traditional strain induced precipitation) and Trial-2 (austenite grain refining by nano-scale precipitate engineering of TiN—NbC in accordance with the present invention) can be found in Table 13, below. The crack tip opening displacement (CTOD) results can be found in Table 15, below. Those skilled in the art will readily appreciate that CTOD can be measured in accordance with ISO/FDIS 3183:2006(E), May 17, 2006 or ANSI/API Specification 5L/ISO 3183. Table 13, below, shows percentage shear area measured in the DWTT as a function of test temperature for Trial-1 and Trial-2.

TABLE 13 Test Temperature, ° C. 0 −15 −30 −40 −50 −60 % Shear area on Trial-1 97 94 90 76 51 50 fracture surface Trial-2 97 94 94 94 87 85

In Trial-1, roughing was carried out in high temperature window in which there is no thermodynamic potential for precipitation of NbC and hence strain induced precipitation of NbC was promoted during finish rolling by pancaking coarse grained austenite. The microstructure exhibited quasi-polygonal ferrite along with acicular ferrite. There is ultra-fine precipitation of NbC in the matrix and pronounced cleavage texture was obtained. The microstructure exhibited low CTOD toughness, characteristic of low crack arrestability.

In comparison, Trial-2 involved lowering the temperature window of roughing (<1040° C.) so that NbC grows on pre-existing TiN to give TiN—NbC composite that could limit grain coarsening of austenite by pinning pressure exerted by TiN—NbC composite. Even though the chemistry of plates processed in Trial 1 and 2 are comparable, the difference lies in time-temperature-deformation schedule in roughing. In Trial-1, roughing was carried out in high temperature window (>1100° C.) to refine the austenite grains below 30 microns. But the finish rolling was applied in low temperature window to promote Tnr by strain induced nucleation of NbC at large undercoolings (<920° C.). There was no protection against grain coarsening of austenite between the end of roughing (1060° C.) and the start of finish rolling (920° C.) and hence austenite grains coarsen to 60-80 microns. In Trial-2, after the first stage of roughing, a second stage of roughing including two heavy passes (>25% reduction) were applied after holding for temperature homogenization in roughing so that Tnr is promoted before the end of roughing by strain-induced growth of NbC on pre-existing TiN, which is well dispersed. By promoting Tnr at about 1040° C., the occurrence of undesirable rotated cube texture is suppressed in Trial-2, whereas high intensity of rotated cube texture is obtained in Trial-1. By promoting strain induced growth of NbC in upstream processing, delayed strain induced precipitation of fine NbC is suppressed in Trial-2. In contrast, the matrix from Trial-1 exhibits high number density of fine NbC precipitates in the matrix. Table 14, below, summarizes the differences in structure and properties between Trial-1 (processed according to the prior art) and Trial-2 (processed according to embodiments of the present invention).

TABLE 14 Trial-1 without nano-scale Trial-2 with nano-scale TiN—NbC precipitate engineering - (HTP) composite precipitate engineering Course austenite grain size before Fine austenite grain size before pancaking pancaking Strain-induced precipitation No strain-induced precipitation High intensity of unfavorable Low intensity of unfavorable {011}<110> texture {011}<110> texture Diffusional transformation product Diffusion-less transformation product of quasi-polygonal ferrite of acicular ferrite Course localized carbide Finely dispersed MA product

By pancaking grain refined austenite in Trial-2, acicular ferrite microstructure was obtained. Both cleavage texture and tendency for ultra-fine precipitation of NbC in the matrix were suppressed. The sample from Trial-2 exhibited excellent DWTT performance and CTOD toughness. Table 15, below, shows the CTOD test results of K60-E2 plate of Trial-1 and Trial-2. Table 16, below, shows the tensile properties of the 32 mm thick gage K60-E2 plate for Trial-1 and Trial-2. This validates the concept of nano-scale precipitate engineering to obtain excellent strength and crack arrestability in thicker gage (32 mm) plate.

TABLE 15 −30° C. −40° C. Trial-1 Trial-2 Trial-2 δμ 0.440 0.667 0.698 2.265 1.933 2.105 (mm) δm 2.119 2.207 1.852 (mm)

TABLE 16 Yield Tensile Strength Strength Yield to Total Area Uniform Trial Rt.0.5(MPa) Rm(MPa) tensile ratio elongation reduction elongation Trial-1 566 681 0.83 23 74 7 Trial-2 583 663 0.88 24 77 7

Control of Percentage Shear Area in DWTT:

By refining the austenite grain size, 100% shear area at low temperature is obtained in 16.4 mm strip in X-90, as shown in FIG. 26. The improvement in DWTT properties is related to suppression of rotated cube texture and promotion of crystallographic variants that suppress brittle fracture in nano-scale precipitate engineering of TiN—NbC composite in plate rolling of 32 mm gage K-60 and strip rolling of 16.4 mm X-90. By refining the austenite grain size, the pancaked austenite thickness is decreased to 6 microns, which acts as the upper bound for domain size. By refining the domain size, the competition from brittle fracture is suppressed, thereby increasing the percentage shear in DWTT at low temperature. FIG. 27 is a comprehensive summary of results that compares the results of the prior art based on strain induced precipitation of NbC in 32 mm K-60 plate rolling with the results from the nano-scale precipitate engineering of 32 mm K-60 of comparable chemistry and 16.4 mm gage X-90 strip in accordance with embodiments of the present invention. As can be seen in FIGS. 26-27, embodiments of the present invention achieve remarkable 100% shear at −60C is attributed to refinement of pancaked austenite thickness and corresponding refinement of crystallographic domain size.

The application of nano-scale TiN—NbC composite precipitation engineering offers a generic platform for austenite grain size control in upstream processing. A potential application to in-line strip rolling involving three roughing stands to produce X-80 grade strip of 15 mm gage is illustrated in FIG. 14 along with critical processing parameters. The suppression of delayed strain induced precipitation of NbC by control of interparticle spacing of TiN—NbC composite is identified as a critical processing parameter to obtain consistent properties in near net shape processing using thin slab casting.

The foregoing examples distinguish the principal differences in processing higher niobium steels between the prior art without intentional austenite grain size control upstream and the present invention based on austenite grain size control and the consequences thereof on product in terms of gage thickness and properties. The salient points are summarized in Table 12.

TABLE 12 ID Prior Art Present Invention 1 No specific nitrogen target N control (nitrogen 0.003-0.004, titanium 0.012- 0.015) 2 Coarse and non-uniformly dispersed TiN Fine uniformly dispersed and high number density of TiN with interparticle spacing of 200- 300 nm 3 Roughing in temperature range where there is First part of the Roughing in temperature range no significant thermodynamic potential for where there is no thermodynamic potential for precipitation of NbC and austenite grains are precipitation of NbC to refine austenite grains by refined by static recrystallization static recrystallization and after cooling to low temperature where significant potential for thermodynamic potential for precipitation of NbC, one or more heavy passes are applied in the second part of roughing to promote strain induced growth of NbC on pre-existing TiN particles to promote adequate Zener drag of TiN—NbC composite to arrest recrystallization and grain coarsening 4 NbC growth on coarse TiN precipitates before High number density of TiN and adequate entry to finish rolling; Inadequate Zener drag to volume fraction of TiN—NbC composite prevent grain coarsening of fine austenite precipitates to give adequate Zener drag to pin grains between the end of roughing and the fine austenite grains from coarsening during start of finish rolling. Potential for delayed cooling between the end of roughing and the start strain induced precipitation of NbC in finish of finish rolling rolling 5 Tnr is by strain induced nucleation of NbC, Tnr by strain induced growth of NbC on pre- which occurs below 960° C. during finish existing TiN occurs at negligible undercoolings, rolling; Recrystallized austenite grains which raise Tnr to 1040° C. in the temperature occurring in the center of thick section at window of roughing. Recrystallized austenite temperature higher than 960° C. will give grains occurring in the center of thick sections, undesirable rotated cube texture upon which is about 100° C. higher than the surface can transformation be pancaked to give brass or copper texture, which upon transformation suppresses the undesirable rotated cube texture 6 Coarsened austenite grains 50-70 μm at entry to Zener limiting austenite grain size (30 μm) at finish rolling Heavy pancaking is required entry to finish rolling; Less pancaking (total (total reduction 66-80%) to achieve target Sv reduction is adequate to achieve target Sv factor, factor. This limits the gage thickness to 16 mm which enables thicker gage to be produced. 7 Lack of control of interparticle spacing in Delayed strain induced precipitation of NbC is upstream processing promotes delayed strain suppressed by control of interparticle spacing of induced precipitation of NbC downstream. TiN—NbC composite; Adequate solute niobium is This renders it difficult to control solute retained in the matrix that promotes acicular niobium required in the matrix to suppress ferrite morphology transformed at low polygonal ferrite fraction in the final temperature with high density and dispersion of microstructure. high angle boundaries that suppress brittle fracture 8 Production limited to thinner gage Production of thicker gage high grade product product (10-17 mm); potential for (17-40 mm) unfavorable texture development less texture related anisotropy

Process Steps for Controlling Austenite Grain Size in Upstream Processing of Austenite:

According to the embodiments of the present invention, the process steps for controlling austenite grain size upstream before entry to finish rolling to produce thicker gage product are given below:

(i) Lower the nitrogen content in the base chemistry to 30-40 ppm and add titanium to the stoichiometric requirement (0.012-0.015 wt percent titanium) to combine with all nitrogen to form in austenite high number density of TiN precipitates in the size range of 10-20 nm with an interparticle spacing of 200-300 nm, before the start of roughing;

(ii) Refine austenite grain size by static recrystallization in rough rolling to a target grain size of 10-30 microns but preferably 10-20 microns by roughing passes in the temperature window above 1080° C. in which there is no significant thermodynamic potential for precipitation of NbC in a base chemistry containing carbon of about 0.045 wt % and niobium content of about 0.09 wt %. This constitutes the first phase of rough rolling aimed at refining the austenite grain size by repeated static recrystallization occurring between pass reductions.

(iii) After the first stage of roughing to refine austenite grain size, hold for about 100-120 seconds so that temperature homogenization occurs through the section thickness and the temperature at the center of the section is lowered to 1040° C.; the corresponding surface temperature will be considerably lower, depending on the section thickness. The second phase of roughing, one or more heavy passes are given to promote strain induced growth of NbC on pre-existing TiN to achieve a Zener pinning pressure of about 0.1 MPa to stop recrystallization and grain coarsening. One heavy pass reduction was used in the second stage of rough rolling in producing 20 mm coil in conventional hot strip mill. More than one heavy pass reduction was used in the second stage of roughing in producing 32 mm plate by nano-scale precipitate engineering of TiN—NbC composite.

(iv) Apply one or more heavy pass reductions (>25%) when the temperature in the center of the section is cooled to about 1040° C. so that there is significant thermodynamic potential for precipitation of NbC to occur. Large strain rolling deformation is applied in a temperature window at which significant thermodynamic potential for precipitation of NbC occurs. This is a critical process step in order to promote strain induced growth of NbC on the pre-existing TiN to target TiN—NbC composites to achieve a Zener pinning pressure of about 0.1 MPa on austenite grain boundary to stop recrystallization during the second stage of roughing. This is achieved by strain induced growth of TiN—NbC composites having a mean size of 25-30 nm with a mean interparticle spacing of 200-300 nm, which can pin austenite of 30 microns grain size in the center of the transfer bar. Since the temperatures at the quarter point and on the surface are lower than the center, there is adequate thermodynamic potential for giving adequate Zener pinning pressure to stop the recrystallization and prevent grain coarsening at these regions. It should be noted that even if high number density of TiN particles with short interspacings are obtained, strain induced growth of NbC will not be promoted unless large strain deformation is coupled to temperature of deformation at which significant thermodynamic potential for precipitation of NbC occurs and therefore Tnr by TiN—NbC will not be promoted. Thus the temperature and pass reduction in second stage of roughing should be coupled to thermodynamic potential for precipitation of NbC.

(v) Apply rapid cooling between the end of roughing and the start of finish rolling so that the temperature of the center of the transfer bar on entry to finish rolling is below 920° C., and (ii) adequate solute niobium >0.03 wt percent, but preferably 0.04 to 0.05 wt percent is retained for strain accumulation during finish rolling and transformation hardening on subsequent accelerated cooling; and

(vi) Control fine austenite grain of about 30 micron size in the transfer bar enables thicker gage (17-40 mm) to be produced with less pancaking in finish rolling. By comparison, heavy pancaking will be required for coarse austenite grain sizes of about 60 microns in conventional thermo-mechanical rolling of higher niobium grades, which results in traditional thinner gage steel products.

Advantages of Embodiments of the Present Invention Based on Nano-Scale TiN—NbC Composite Precipitate Engineering for Austenite Grain Size Control:

The foregoing examples are given to demonstrate how nano-scale precipitation engineering of TiN—NbC composite precipitates can be used for austenite grain size control in upstream processing of austenite to derive benefits in (i) producing thicker gage product (>17 mm) with excellent strength and fracture toughness at low temperature as measured by DBTT and DWTT, (ii) obtaining more uniform microstructures, and (iii) minimizing unfavorable crystallographic texture related problems. Strain induced growth of NbC on pre-existing TiN to form TiN—NbC composites at low undercoolings is a paradigm shift from conventional technology based on strain induced nucleation of NbC at large undercoolings. The increase in Tnr by TiN—NbC composites is about 100° C. more than Tnr by NbC. This enables Tnr to be obtained in rough rolling, which prevents grain coarsening between the end of roughing and the start of finish rolling. By raising Tnr, the center of thick sections can be pancaked before transformation, thereby suppressing undesirable rotated cube texture from occurring in the center of thicker gage product. By controlling the interparticle spacing of TiN—NbC composites, it is shown that delayed strain induced precipitation of NbC can be suppressed, thereby retaining adequate solute niobium for hardenability control to produce fine acicular ferrite with crystallographic variant control to suppress brittle fracture. Nano-scale precipitate engineering enables control of undesirable rotated cube texture and promotes desirable crystallographic variants that suppress brittle fracture. Thus the target structure achieved by nano-scale precipitate engineering of TiN—NbC composites is acicular ferrite of 1-3 microns transformed at low temperature by shear mechanism, which is characterized by high density and dispersion of high angle boundaries that suppress brittle fracture. The target structure exhibits minimal rotated cube texture and minimal ultra-fine precipitates that contribute to strengthening. The increase in strength by precipitation hardening will increase the ductile to brittle transition temperature (DBTT) and hence it is not a desirable mechanism to obtain toughness at low temperature. Impact transition temperature will be raised by 0.26° C. per 1 MPa increase in yield strength. Nano-scale precipitate engineering of TiN—NbC composites is designed to improve crack arrestability of running crack to enhance public safety. The design is aimed at promoting (i) morphological microstructural parameters of fine acicular ferrite grain size of 1-3 microns with good matrix plasticity to increase resistance to ductile fracture propagation, and (ii) crystallographic parameters to prevent brittle fracture by (a) suppressing undesirable rotated cube texture that promotes brittle fracture, and (b) promoting crystallographic high angle boundaries to arrest micro-cracks before they grow to initiate brittle fracture. The roles of TiN—NbC and solute niobium in microstructural engineering of target structure for enhanced crack arrestability are summarized in FIGS. 28 and 29 respectively.

These examples are for illustrative purposes only and the invention is not intended to be limited to any of the specific examples. However, it will be understood by those skilled in the art that modifications and changes may be made to the present invention to combine other elements having a high affinity for nitrogen and carbon similar to titanium and niobium without departing from their scope of controlling particle interspacing and size independently to bring about adequate pinning pressure on the austenite boundary and prevent austenite grain coarsening upstream. 

What is claimed is:
 1. A process for controlling austenite grain size in austenite processing through nano-scale precipitate engineering of TiN—NbC composites to produce thicker gage product of niobium microalloyed steel, comprising: (i) controlling the base chemical composition of a steel product to include Amount Element (wt %) N 0.003-0.004 Ti 0.012-0.015 C 0.03-0.07 Nb 0.07-0.15

wherein controlling the nitrogen to 0.003-0.004 wt % and controlling the Ti to 0.012-0.015 wt % in the steel product before rolling produces TiN precipitates with a mean interparticle spacing in the range of about 200-300 nm and a mean size in the range of about 10-20 nm; (ii) conducting a first stage of roughing on the steel product with rolling passes in a temperature window ranging from 1080° C. to 1150° C. to promote refinement of austenite grains from 300 microns to less than 30 microns by repeated static recrystallization between the rolling passes; (iii) cooling the steel product for 100-120 seconds to allow for temperature homogenization and for the center of the steel product to cool to a temperature of about 1040° C.; (iv) conducting a second stage of roughing on the steel product, wherein the second stage of roughing includes conducting at least one heavy rough rolling pass reduction (>25%) wherein the center of the steel product is at about 1040° C. in order to promote strain induced growth of NbC on pre-existing TiN to form TiN—NbC composites having a mean size of 25-30 nm with a mean interparticle spacing of 200-300 nm pre-set by the interparticle spacing of the TiN to stop the static recrystallization in the center of a thick section of the steel product by pinning the austenite grain boundaries with a Zener pinning pressure of about 0.1 MPa, raising the temperature of no recrystallization in the roughing window to counteract the driving force of recrystallization and prevent grain coarsening of austenite grains beyond 30 microns in the steel product; (v) cooling the center of the steel product to a temperature at or below about 920° C. to retain solute niobium (Nb>0.03 wt %) in the matrix for transformation hardening; (vi) conducting finish rolling on the steel product to achieve a total rolling reduction of 70% to pancake the austenite grains of about 30 microns to target a pancaked austenite thickness of less than 10 microns; and (vii) applying accelerated cooling to the steel product after finish rolling and its austenite grains having a large surface to volume ratio (Sv factor), large strain accumulation and solute niobium (>0.03 wt % Nb) in the matrix to produce 1-3 micron size acicular ferrite grains upon transformation to obtain high strength and toughness at low temperature.
 2. A process as recited in claim 1, wherein the niobium in the base chemical composition increases the hardenability in order to promote acicular ferrite transformation at low temperature by shear transformation to promote high angle boundaries to arrest microcracks, thereby suppressing brittle fracture.
 3. A process as recited in claim 1, wherein the second stage of roughing aids in strain induced growth of NbC on pre-existing TiN by accelerated diffusion due to dislocations generated by deformation from the rough rolling to form the TiN—NbC composite.
 4. A process as recited in claim 1, wherein the TiN—NbC composite size of 25-30 nm and interparticle spacing of 200-300 nm results in an austenite grain size ranging from 20-40 microns before conducting finish rolling, thereby increasing the surface to volume ratio of austenite grains (Sv factor) that enables production of thicker gage steel product.
 5. A process as recited in claim 1, wherein the TiN—NbC interparticle spacing of 200-300 nm is configured to suppress delayed strain induced precipitation of NbC, which renders it difficult to control solute niobium required in the matrix to control the transformation structure in downstream processing.
 6. A process as recited in claim 1, wherein the Zener pinning pressure of about 0.1 MPa results in austenite grains in the center of the steel product, which is at a higher temperature than the surface of the steel product by about 100° C., being fully pancaked to modify the cubic texture of recrystallized austenite grains to brass or copper texture before phase transformation in order to suppress the formation of undesirable rotated cube cleavage texture upon transformation, thereby suppressing brittle fracture in the final product.
 7. A process as recited in claim 1, wherein the dissolution temperature of NbC is more than 1040° C.
 8. A process as recited in claim 1, further comprising processing the steel product by at least one of conventional plate rolling, conventional hot strip rolling, steckel mill rolling, or near net shape processing.
 9. A process as recited in claim 1, wherein the steel product is at least one of line pipe steel, infra-structure steel, or super-martensitic stainless steel.
 10. A process as recited in claim 1, wherein the crystallographic texture-related anisotropic properties of the resulting steel product are minimized.
 11. A process as recited in claim 1, further comprising substituting titanium partially or fully in the base chemistry with a member of the group consisting of Zr, Hf, Ta, W, V, Cr, Mo, Al and mixtures thereof, each with high affinity for nitrogen to form nano-scale precipitates on which NbC can grow epitaxially to give composite precipitates.
 12. A process as recited in claim 1, further comprising partially substituting niobium in the base chemistry with other microalloying elements with high affinity for carbon selected from the group consisting of Zr, Hf, Ta, W, V, Cr, Mo, and mixtures thereof, each to give composite precipitates.
 13. A process as recited in claim 1, further comprising substituting solute niobium on entry to finish rolling with other elements, which exhibit solute drag comparable to niobium.
 14. A process as recited in claim 1, wherein the steel product exhibits a gage thickness of about 17-40 mm.
 15. A process for controlling austenite grain size in austenite processing through nano-scale precipitate engineering of TiN—NbC composites to produce thicker gage product of niobium microalloyed steel, comprising: (i) controlling the base chemical composition of a steel product to include Amount Element (wt %) N 0.003-0.004 Ti 0.012-0.015 C 0.03-0.07 Nb 0.07-0.15

(ii) conducting a first stage of roughing on the steel product with rolling passes in a temperature window ranging from 1080° C. to 1150° C. to promote refinement of austenite grains from 300 microns to less than 30 microns by repeated static recrystallization between the rolling passes; (iii) cooling the steel product for 100-120 seconds to allow for temperature homogenization and for the center of the steel product to cool to a temperature of about 1040° C.; (iv) conducting a second stage of roughing on the steel product, wherein the second stage of roughing includes conducting at least one heavy rough rolling pass reduction (>25%) wherein the center of the steel product is at about 1040° C.; (v) cooling the center of the steel product to a temperature at or below about 920° C. to retain solute niobium (Nb>0.03 wt %) in the matrix for transformation hardening; (vi) conducting finish rolling on the steel product to achieve a total rolling reduction of 70% to pancake the austenite grains of about 30 microns to target a pancaked austenite thickness of less than 10 microns; and (vii) applying accelerated cooling to the steel product after finish rolling to produce 1-3 micron size acicular ferrite grains upon transformation to obtain high strength and toughness at low temperature. 